颗粒增强型铝基复合材料电阻点焊焊接工艺 外文翻译.pdf

Science Press Available online at Trans Nonferrous Met SOC China 16 2006 818 823 Transactions of Nonferrous Metals Society of China bi Electron beam welding of SiCp LD2 composite CHEN Mao ai I WU Chuan song i j j ZOU Zeng da P Key Laboratory of Liquid Structure and Heredity of Materials Ministry of Education Shandong University Ji nan 250061 China Received 18 October 2005 accepted 23 November 2005 Abstract The 2 mm thick SiC electron beam welding interfacial reaction 1 Introduction Sic particle reinforced aluminum composites SiCdAl possess excellent properties e g high strength to mass ratio high stiffness to mass ratio good wear resistance and higher service temperature therefore they are considered a kind of promising structural material With the gradual reduction in production cost an increase in uses and application may be expected in the next years For this reason welding technologies of this lund of material have attracted a lot of attention recently Previous studies reveal that almost all fusion welding processes produce deleterious microstructural alterations when they are applied to SiCdAl composites l 91 In TIG and MIG welding pool the strong rejection of Sic particles by solidification front results in microsegregation and inhomogeneous particle distribu tion thus reducing the reinforcement efficiency 1 31 The high hydrogen content in the A1 matrix composites especially the one produced by powder metallurgical routes and low flowability of molten pool tend to result in serious porosity in both weld and HAZ 2 51 The vacuum degassing prior to welding can eliminate the porosities in the TIG and MIG welds 2 3 5 During welding serious interface reaction between Sic particles and A1 matrix occurs and plate like A14C3 precipitates form which are very brittle and deleterious to the pro perties of the weld 2 3 7 8 However through applying Al Si filler metal and or lowering heat input the interfacial reaction can be reduced or suppressed 2 3 By using solid state joining the above mentioned problems can be eliminated Therefore much more research work has been done in solid state joining such as diffision bonding and friction welding of Sic particle reinforced aluminum composites and results are encouraging 10 121 However the low production efficiency and suitability limit the application fields of those procedures Thus it is still necessary to reconsider the possibility of application of the more productive welding processes fusion welding In this paper the influence of heat input on microstructure and properties of the EB electron beam welds in SiCdLD2 were studied The emphasis was placed on the analysis of effect of welding cycle on the particle distribution A1 Sic interface morphologies and the interfacial reaction products in the welds The influence of interfacial reaction on the properties and fracture mechanism of the EB welds were also discussed The mechanism of interfacial reaction was also reviewed Foundation item Project 2004BS05010 supported by Shandong Youth Doctoral Foundation China Corresponding author CHEN Mao ai Tel 86 53 1 8839271 1 E mail chenmaoai sdu CHEN Mao ai et aL Trans Nonferrous Met SOC China 16 2006 819 2 Experimental 2 1 Material The material used in this investigation was 2 mm thick plate of Sic particle reinforced LD2 aluminum SiCdLD2 The volume fraction of Sic particle was 10 2 2 Welding machine and procedure The 2 nun thick as received plate was cut into to be welded pieces with 500 mm in length and 50 mm in width Before welding the pieces were cleaned with acetone thoroughly to eliminate dirt and grease Welding experiments were performed with EBW2100115 15OCNC The welding parameters are listed in Table 1 Heat inputs in Table 1 were calculated as the product of accelerating voltage and beam current divided by traveling speed 2 3 Microstructure analysis and mechanical testing Metallurgical specimens were sectioned from the weldment in transverse to the welding direction The specimens were first polished with diamond paste and water and finally polished with diamond paste and alcohol to prevent A14C3 that may be present in the weld from decomposing After polishing the specimens were cleaned with acetone Prepared specimens were observed with Neophto 20 optical microscope Quantitative metallography was used to measure the mean size of A 2 formation of A14C3 near the interface The dissolution of the interior part of the Sic particles results from the crystal structure defects TEM analysis also reveals that needle like A14C3 precipitates observed under optical microscope consist of a lot of discrete A14C3 tablets Fig 3 a These tabular A14C3 precipitates have different crystal orientations Fig 3 b although they nearly have the same geometric direction The interfacial reaction products between Sic and A1 include A4C3 and free Si but no Si crystals are found at the Sic A1 interface While equiaxed Si crystals in single or a group are observed in the A1 region Fig 4 a Other researchers 5 6 also did not found Si on the Sic surface while studying the interfacial products of SiCp Al composites with TEM or HREM Those observations suggest that Si produced by interfacial reaction have diffused into the A1 matrix at a high rate In addition Al Fe Si phases are found on the grain boundary ofAl in the EB weld Fig 4 b Fig 5 shows the particle distributions in EB welds and as received SiC LD2 composite The particles in the Fig 1 Typical microstructures of EB welds in SiC LD2 made at different heat inputs a 30 J mm b 36 J mm c 42 J m d 48 J llUn Fig3 TEM images of Sic A1 interface in 30 J mm fusion zone a AI4C3 tablets gowing into A1 region b A14C3 tablets growing into Sic particle CHEN Mao ai et al Trans Nonferrous Met SOC China 16 2006 82 1 Fig3 TEM morphologies of needle like A14C3 in 42 J mm fusion zone and its diffraction pattern a TEM morphology of needle like A b Diffraction pattern Fig 4 Si crystals a and Al Fe Si phase b in A1 region as received composite exhibit a slight laminar distribu tion along the rolling direction Fig S a And the Sic particles in the EB weld are distributed much more uniformly Figs S b and 5 c The uniform distribution of particle in weld may be attributed to the key hole effect and the rapid cooling rate of the EB pool During welding the vaporizing of A1 matrix and colliding of the atom to the composite result in a formation of keyhole The melted A1 around the keyhole is subjected to a violent stirring movement which might make the Sic particles distribute uniformly in the weld pool The particle distribution in weld depends on the particle distribution in molten pool and the rejection of the solidification It is generally accepted that solidification rate influences the rejection effect of solidification front to the particle As the solidification rate increases the rejection effect decreases 151 According to the estimation of Loloyd the cooling rate of the EB welds is in the range of approximate 3 000 6 000 C s at a heat input of 33 5 J mm l6 With such a high cooling rate almost no particle rejections exist as a result the particle distributions in the EB fusion zones are fairly uniform Hydrogen porosities often observed in TTG welds and MIG welds are not found in the fusion zone of the EB welds The possible reason may be as follows the hydrogen in the molten pool can escape easily because of the vacuum environment of EB welding and the violent stirring of the molten pool furthermore no cracks are observed in the EB welds due to the rapid cooling of molten pool When welded at a heat input of 30 J m the work pieces were not completely penetrated therefore no mechanical property tests were carried out for the 30 Jlmm weld All other weldments were tested under as welded condition and failed in the weld metal during testing The maximum tensile strength measured is about 60 of the as received composite strength the maximum elongation is about 40 of the as received composite elongation This indicates that the interfacial reaction exerts more effect on the ductility than on the tensile strength The tensile testing data for individual EB weld are summarized in Table 2 It can be seen from Table 2 that the tensile strength and ductility elongation decrease with the heat input increasing During tensile fracture of SiCdAl composite the crack propagation path involves Sic particles A1 matrix and particle matrix interface 171 SEM examination of different fracture surfaces shows that the extent of ductile dimpling decreases with the heat input increasing and the failure mechanism varies with the heat input Fig 6 shows 822 CHEN Mao ai et aVTrans Nonferrous Met SOC China 16 2006 Fig 5 Particle distributions in EB welds and as received composite a As received composite b 30 J mm EB weld c 42 J mm EB weld Table 2 Tensile ProDerties of EB welds on SiCJLD2 Elongation Remark Heat input1 strength Serial No J mm MPa 303 5 6 T6 As received O composite Incomplete 1 30 penetration 2 36 204 2 1 As welded 198 2 3 196 2 6 3 42 169 1 3 As welded 153 1 4 160 1 6 4 48 120 As welded 112 107 Fig 6 SEM fractographs of EB welds metal and as received composite a As received composite b 36 Jimm weld c 48 J mm weld the typical fracture morphologies of as received composite and individual EB welds The fracture surface of the as received composite exhibits typical ductile morphology with a lot of fine dimples in the A1 matrix region between particles there are some fractured particles on the fracture surface Fig 6 a This indicates that the fracture process of as received material is controlled mainly by matrix ductile failure and particle fracture mechanism The fracture surface of the 36 J mm weld also exhibits evident ductile morphology the quantity of dimples is much less than that of as received composite Although most particles survive the welding cycle few fractured particles are observed on the fracture surface Fig 6 b This is because that the A14C3 tabular and the dissolution of the Sic surface during welding reduce the strength of Sic A1 interface The cracks progress along the Sic A1 interface when they meet the particles In this CHEN Mao ai et al Trans Nonferrous Met SOC China 16 2006 823 case the weld fracture process is controlled mainly by matrix ductile failure and interface fracture mechanism The energy required by the crack propagation along the weak interface is much low Therefore the strength and elongation of the weld are obviously lower than that of the as received composite Table 2 although the particle in the weld is distributed much more uniformly than in the as received composite The fracture surfaces of the 42 J mm and 48 J mm welds present typical brittle morphologies in the A1 matrix region and no fractured Sic particles are observed Fig 6 c The corresponding strengths and elongations of the welds are very low Table 2 The serious interfacial reaction is responsible for the loss of strength and elongation The formation of network of acicular AI4C3 which is very brittle and susceptible to initiate greatly degrades the properties of the weld Furthermore another interfacial product Si results in brittle aggregates such as Si crystal Al Fe Si intermetallic compound and Al Cu Si eutectic on the grain boundary of the A1 matrix Those brittle aggregates degrade the ductility and strength of A1 matrix further 4 Conclusions 1 At all the heat inputs used interfacial reaction occurs and A4C3 precipitates forms Both A tablets growing from interface into the A1 region and those growing into the Sic particle are observed This indicates that A14C3 precipitates can form only after dissolving of Sic in liquid Al 2 The size and the amount of AI4C3 precipitates decrease with the heat input decreasing When the 2 mm thick SiCdLD2 palate is welded at a heat input of 30 J mm no needle like A14C3 precipitates form Sic particle distribution in EB is more uniform than that in as received composites When the heat input increases to 36 Jimm a few small sized needle like precipitates are observed in the top center of the fusion zone and sound EB welds can be obtained 3 The needle like A C3 precipitate observed under optical microscope consists of many tabular A C3 crystals which have different orientations 4 With the heat input increasing the fracture mechanism changes from ductile one to brittle one the quantity of fractured particles on the fracture face decreases and the strength and ductility of the weld decrease The degradation of the properties of the weld are mainly caused by the formation of acicular A14C3 and Si rich phase References MIDDLING 0 T GRONG 0 Joining of particle reinforced AI Sic MMCs J Key Engineering Materials 1995 104 107 2 355 372 CHEN Mao ai WU Chuan song GAO Jian guo Welding of Sic particle reinforced 6061 Al matrix composite with P TIG J Trans Nonferrous Met SOC China 2002 12 5 805 808 CHEN Mao ai WU Chuan song WANG Jim guo Effect of composition of welding wire on microstructure and mechanical properties of weld metal in Sic particle reinforced 6061A1 matrix composite J Transactions of the China Welding Institution 2003 24 5 69 72 in Chinese DEVLETIAN J H SiCiAl metal matrix composite welding by a capacitor discharge process J Welding Journal 1987 66 6 33 39 AHEARN J S COOKE C FISHMAN S G Fusion welding of Sic reinforced A1 composite J Metal Construction 1982 14 4 192 197 NIU Ji tai WANG Mu zhen LAI Zhong hong Mechanism of laser welding for SiCJ6061AI composite J Transactions of the China Welding Institution 2000 21 1 1 4 in Chinese URENA A ESCALERA M D GIL L Influence of interface reaction on kacture mechanisms in TIG arc welded aluminum matrix composites J Composites Science and Technology 2000 20 4 613 622 NARENDRA B DAHOTRE T DWAYNE M MCCAY M H Laser processing of a SiC A1 alloy metal matrix composite J Journal of Applied